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Microstructure and fracture of alloyed austempered ductile iron
Materials Characterization 57 (2006) 211 – 217
Microstructure and fracture of alloyed austempered ductile iron Olivera Erić a , Dragan Rajnović b , Slavica Zec a , Leposava Sidjanin b , Milan T. Jovanović a,⁎ Institute of Nuclear Sciences “Vinča”, Belgrade, Serbia and Montenegro University of Novi Sad, Faculty of Technical Sciences, Novi Sad, Serbia and Montenegro a
b
Received 13 July 2005; accepted 6 January 2006
Abstract An investigation has been conducted on two austempered ductile irons alloyed with Cu and Cu + Ni, austenitized at 900°C and austempered at 350°C. The microstructure and fracture mode developed through these treatments have been identified by means of light and scanning electron microscopy and X-ray diffraction analysis. Impact energy measurements were performed on un-notched Charpy specimens. The maximum value of retained austenite volume fraction observed in the material alloyed with Cu + Ni was higher than in that alloyed with Cu austenitized and austempered under the same conditions. This led to the material alloyed with Cu + Ni having higher impact energy and substantial plastic deformation. © 2006 Elsevier Inc. All rights reserved. Keywords: Alloyed ductile iron; Bainitic transformation; Retained austenite; Impact energy
1. Introduction In recent years, there has been a significant interest in the processing and development of austempered ductile cast irons (ADI). Heat-treated ADI have been reported to have excellent mechanical properties and it appears that this material can be processed into major engineering components with a wide range of versatile properties [1–5]. ADI is a heat-treated alloyed nodular cast iron with a microstructure that consists of high carbon austenite and acicular bainitic ferrite with graphite nodules dispersed in the matrix. When discussing the microstructure of ADI, it is necessary, according to some authors [6], to ⁎ Corresponding author. Tel.: +381 11 458 222; fax: +381 11 3440 100. E-mail address: [email protected] (M.T. Jovanović). 1044-5803/$ - see front matter © 2006 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2006.01.014
distinguish the residual austenite that exists at the isothermal transformation temperature from the retained austenite that remains untransformed at ambient temperature. During the isothermal transformation, carbon partitions between ferrite laths and austenite. Undesirable phases, such as martensite and iron carbides, may also be present in smaller quantities, but it is understood that a maximum retained austenite volume fraction is important if good mechanical properties are to be achieved. The base iron chemistry and alloy additions to ductile iron play important roles in ADI technology. The addition of alloying elements during the production of ADI is often considerably higher than the levels used in the production of conventional grades of ductile irons [7]. The alloy in the austenitized condition (obtained by holding at a temperature in the range 820–950 °C long enough to dissolve the austenite with up to 2.1 wt.% carbon) is quenched to the austempering temperature
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and subjected to isothermal transformation at temperatures between 250 and 450°C before cooling to ambient temperature. The microstructure can be controlled through the variation of alloying elements, temperature and the time of austempering. The choice of austempering parameters is critical. The isothermal heat treatment should be terminated at the point that allows an optimum balance between bainitic ferrite and retained austenite, but before carbide precipitation. The formation of carbides may cause the structure to become embrittled and thus degrade the mechanical properties of ADI [5]. Since the majority of ADI components should possess satisfactory austemperability, alloying elements serve to delay the transformation of austenite [8–13]. Due to a severe restriction on the maximum thickness of bulk unalloyed ADI that can be heat-treated successfully, alloying with Si, Mn, Mo, Ni and Cu is aimed at providing a desired hardening ability. The importance of alloying additions depends on their relative effectiveness on the austempering process. Si has a very important role because the addition of more than 2 wt. % Si may retard carbide precipitation, producing carbide-free austenite in a bainitic microstructure. The effect of Si and of individual or combined additions of Cu, Mo, Ni, and Mn on the transformation characteristics of ADI have been reported earlier [9,10]. Mn and Mo delay the austempering process, but Cu does not affect the carbon diffusion in austenite or the stability of austenite. However, it has been reported that Cu suppresses carbide formation in lower bainite, whereas Ni, which acts in a similar way to Cu if present in excess of 0.5%, slows down the bainitic reaction, causing the formation of martensite at austenite cell boundaries on cooling [10]. The focus of this paper is the influence of additions of Cu and Cu + Ni on the properties of ADI. The specific items of study are the effects of austempering time and alloy composition on the microstructure, the impact properties and the fracture mechanism after austenitizing at 900°C and subsequent austempering at 350°C.
spectrometer, was as follows: (a) 3.6 C; 2.5 Si; 0.28Mn; 0.04Cr; 0.45Cu; 0.01P; 0.014 S; 0.066 Mg, and (b) 3.07C; 2.15.Si; 0.26Mn; 0.04 Cr; 1.6Cu; 1.5 Ni (the balance in both alloys was Fe). After austenitizing in a protective argon atmosphere at 900 °C for 2h, the specimens were rapidly transferred to a salt bath at the austempering temperature of 350 °C, held there for between 1 and 6 h, then air-cooled to room temperature. Fracture toughness was measured on un-notched Charpy specimens (55 mm × 10 mm × 10 mm) machined from the Y blocks. The austempered Charpy specimens were tested at room temperature in a standard impact testing machine using an applied load of 150 J. At least three specimens were tested for each heat treatment. The fracture surfaces were examined by a JEOL JSM6460LV scanning electron microscope (SEM) operating at 25 kV. Standard metallographic preparation techniques (mechanical grinding and polishing followed by etching in Nital) were applied prior to light microscope (LM) examinations. A Leitz metallographic microscope was used for microstructural characterization, and an
2. Experimental details Two ductile irons were produced in a commercial high frequency electro-induction foundry furnace. The Fe–Si–Mg (40 wt.%, 50 wt.% and 5 wt.%, respectively) alloy was used for the Mg treatment of the molten metal. The inoculant Fe–Si (with 80 wt.% Si) was added to the molten metal immediately before casting. The melt was poured from about 1420 °C into standard 25.4 mm Y block sand molds (ASTM A-395). The chemical composition (in wt.%) of these irons, obtained by
Fig. 1. LM. Microstructure of ADI alloyed with (a) Cu and (b) Cu + Ni. GN—graphite nodule; F—ferrite; P—pearlite.
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analyzed in order to calculate the volume fraction of phases present, i.e.: Iγ =Iα ¼ Rγ =Rα d Xγ =Xα where Iα and Iγ are integrated intensities from a given hkl of α and γ phase, Xα and Xγ are the volume fractions
Fig. 2. LM. Microstructure of ADI alloyed with Cu after austempering at 350°C for (a) 1h, (b) 2h, and (c) 6h. RA—retained austenite; AF— acicular ferrite; M—martensite; AFS—acicular ferrite sheaves.
Opton Axioplan light microscope equipped with the software “Vidas” program was applied to the measurement of the distribution of graphite nodules. Changes in the volume fraction of austenite retained during austempering were determined by the X-ray diffraction technique using a Siemens D-500 diffractometer with Ni-filtered Cu Kα radiation. The diffractograms were
Fig. 3. LM. Microstructure of ADI alloyed with Cu + Ni after austempering at 350°C for (a) 1h, (b) 3h and (c) 6h. RA—retained austenite; AF—acicular ferrite; M—martensite; AFS—acicular ferrite sheaves; MB—mixed bainite; FF—feathery bainitic ferrite.
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of α and γ phase. Rα and Rγ values may be calculated according to the equation
diffraction analysis was also used for identification of the phases present.
R ¼ ð1=m2 Þ½jF 2 jpð1 þ cos2 2hÞ=sin2 hcoshÞe−2M
3. Results and discussion
where F is a structural factor, p multiplicity factor, θ the Bragg angle, ν the volume of the unit cell, (1 + cos22θ) / sin2θcosθ) corresponds to Lorenz polarization factor, and e− 2M is a temperature factor (a function of θ). X-ray
3.1. Microstructure Light micrographs of cast ADI alloyed with Cu and with Cu + Ni are shown in Fig. 1(a) and (b), respectively.
Fig. 4. X-ray diffraction patterns of ADI alloyed with (a) Cu and (b) Cu + Ni, after austempering at 350°C.
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Fig. 5. The variation of retained austenite content with austempering time at 350°C in ADI alloyed with Cu and Cu + Ni.
In the microstructure of the ADI alloyed with Cu, the pearlite, at 70vol.%, outweighs the ferrite (30 vol.%) and the mainly spherical graphite nodules (the nodule count was around 100 nodules/mm2) (Fig. 1(a)). The microstructure of the ADI alloyed with Cu + Ni consists of graphite spheroid nodules (nodule count over 100 nodules/mm2) in a fully pearlitic matrix (Fig. 1(b)). The microstructures of the ADI alloyed with Cu and austempered at 350°C for 1, 2 and 6 h are shown in Fig. 2(a)–(c), respectively. In specimens austempered for 1h the predominant phase is martensite, whereas only small amounts of acicular ferrite (lower bainite), sporadically packed in sheaves, and retained austenite are present (Fig. 2(a)). At short austempering times, the carbon content is insufficient to retain stable austenite, which consequently transforms to martensite during air-cooling to the room temperature. However, after austempering for 2h, the presence of martensite could not be detected and the structure consisted only of acicular bainitic ferrite and retained austenite (Fig. 2(b)). This indicates that carbon enrichment was sufficient to stabilize the austenite even after cooling. During longer austempering (6h), the amount of retained austenite decreases and acicular bainitic ferrite dominates the microstructure. A small fraction of martensite is also present (Fig. 2(c)). In this case, the large amount of bainitic ferrite leaves less austenite available for retention. Also, the formation of bainitic ferrite consumes carbon, rendering the austenite unstable and promoting transformation to martensite. Microstructures of ADI alloyed with Cu + Ni austempered at 350°C for 1, 3 and 6 h are shown in Fig. 3 (a)–(c), respectively. After austempering for 1 h, the microstructure reveals blocky darkened areas containing
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a high fraction of plate-like martensite with a low fraction of austenite (Fig. 3(a)). After 3 h of austempering, there is an appreciable decrease in the amount of martensite and an increase in the amount of acicular bainitic ferrite and retained blocky austenite (Fig. 3(b)), suggesting that a balance between volume fractions of acicular bainite and stable retained austenite might be reached. After austempering for 6 h, only small amounts of martensite and retained austenite are present in the microstructure, and a mixture of acicular and feathery bainitic ferrite (looks like upper bainite) prevails (Fig. 3 (c)). Similar to ADI alloyed with Cu, a decrease in the amount of retained austenite was the consequence of a high volume fraction of acicular and plate bainitic ferrite. Also, there is some evidence of a mixed mode of transformation, i.e. in some areas acicular and plate bainitic ferrite overlap, a phenomenon already reported in the literature [13]. Comparing the austempered microstructures of both ductile irons, one could observe that alloying with Cu + Ni promotes the transformation of upper bainitic ferrite. The LM observations are supported by the X-ray diffraction patterns of ADI alloyed with Cu and Cu + Ni for different austempering times, shown in Fig. 4(a) and (b), respectively. The profile fitting of the 111 and 110 lines of austenite and ferrite, respectively, were identified in nearly all cases. However, doublets, due to the tetragonality of the martensite unit cell, are superimposed on the strong 110 ferrite line and could not be resolved into separate lines because of large broadening. Diffractograms of ADI alloyed with Cu (Fig. 4(a)) show that the intensity maximum of the 111 line was reached after austempering for 2h, whereas the presence of ferrite and martensite dominate after 1 and
Fig. 6. The effect of austempering time on impact energy after austempering at 350 °C of ADI alloyed with Cu and Cu + Ni.
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6h. The situation is similar with ADI alloyed with Cu + Ni, except that the 111 line could not be detected after the shortest austempering time. The presence of a broad 110 line suggests that ferrite and martensite are the main microconstituents (Fig. 4(b)). After austempering for 3 h, the 111 line exhibits an intensity maximum, indicating a high volume fraction of austenite in the microstructure. Fig. 5 shows the variation of retained austenite content with austempering time in both ductile irons. These results show that the trend in variation of the volume fraction of retained austenite with austempering time is similar for both materials. The maximum value of retained austenite (16.5 vol.%) noted in the ADI alloyed with Cu is achieved after 2h. However, with the ADI alloyed with Cu + Ni, the retained austenite reaches a maximum (19 vol.%) after 3 h. It appears that alloying with Cu + Ni delays the transformation kinetics of residual austenite (isothermal) until the maximum of
the volume fraction is reached. These results are in agreement with the LM and X-ray diffraction findings. 3.2. Impact properties The effect of alloying ADI with Cu and Cu + Ni on the impact energy of specimens previously austempered at 350 °C is illustrated in Fig. 6. The results show that the impact energy increases as the austempering time increases up to a maximum value, after which it decreases. The maximum impact energy (110 J) of ADI alloyed with Cu is achieved after austempering for 2 h, whereas alloying with Cu + Ni postpones the maximum (122 J in this case) to 3 h. The fact that the maximum impact energy corresponds to the maximum value of retained austenite in both cases suggests that the impact values depend on the retained austenite content. That is, with increasing content of retained austenite up to the maximum value, the impact energy also increases,
Fig. 7. SEM. Fracture morphology after austempering at 350 °C. ADI alloyed with Cu for (a) 2h and (c) 6h; ADI alloyed with Cu + Ni for (b) 3 h and (d) 6h.
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then, as the retained austenite content decreases, so the impact energy starts to decrease.
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of retained austenite up to a maximum value impact energy increases, then a decrease occurs with the decrease of retained austenite.
3.3. Fractography Acknowledgements The fracture surfaces of un-notched impact specimens after austempering at 350 °C for times corresponding to maximum impact energy (2 and 3h) and to the subsequent lowest impact energy (after 6h) are shown in Fig. 7. The micrographs of specimens with highest values of retained austenite volume fraction and impact energy exhibit a typical ductile fracture mode (Fig. 7(a) and (b)). The somewhat higher number of dimples observed in the Cu + Ni alloy (Fig. 7(b)) is associated with the higher values of toughness and retained austenite volume fraction. After austempering for 6h, the fracture mode is changed. In the ADI alloyed with Cu a mixture of ductile and cleavage fracture prevails (Fig. 7(c)) but in the Cu + Ni alloyed ADI, only cleavage fracture can be seen (Fig. 7(d)). These results are in accordance with the results of the impact energy testing because the Cu alloyed ADI exhibited the higher impact energy after austempering for 6 h. 4. Conclusions The effect of alloying with Cu and Cu + Ni on the microstructure and impact properties of austempered ductile iron has been studied. Addition of Cu + Ni delays the transformation kinetics of residual austenite (isothermal) resulting in a shift of the maximum of volume fraction of retained austenite to 3 h of austempering, compared to 2 h in ADI alloyed with Cu. In the same time, higher maximum value of the volume fraction of retained austenite was observed in ADI alloyed with Cu + Ni. In the same ADI a substantial plastic deformation at the peak of impact energy is associated with the highest volume fraction of retained austenite. It has been demonstrated that the volume fraction of retained austenite strongly affects impact energy of both irons, i.e. with increasing content
This work was financially supported by the Ministry of Science and Environment of the Republic of Serbia through the Project No. 1970. References [1] Eric O, Sidjanin L, Miskovic Z, Zec S, Jovanovic MT. Microstructure and toughness of CuNiMo austempered ductile iron materials. Mater Lett 2004;58:2707–11. [2] Dodd J. High strength, high ductility ductile irons. Mod Cast 1978;68:60–6. [3] Kovacs BV. Austempered ductile iron—fact and fiction. Mod Cast 1990;36:38–41. [4] Moore DJ, Rouns TN, Rundman KB. Structure and mechanical properties of austempered ductile irons. AFS Trans 1985;93:705–18. [5] Kovacs BV, Keough JR. Physical properties and application of austempering grey iron. AFS Trans 1993;101:283–91. [6] Yescas MA, Bhadhesia HKDH. Model for the maximum fraction of retained austenite in austempered ductile iron. Mater Sci Eng A 2002;333A:60–6. [7] Rouns TN, Rundman KB, Moore DM. On the structure and properties of austempered ductile cast iron. AFS Trans 1984;90:815–40. [8] Yu KS, Loper Jr CK, Cornell HH. The effect of molybdenum, copper and nickel on the microstructure, hardness and hardenability of ductile cast iron. AFS Trans 1986;94:255–64. [9] Erić O, Rajnović D, Šidjanin L, Zec S, Jovanović TM. An austempering study of ductile iron alloyed with copper. J Serb Chem Soc 2005;70(7):1015–22. [10] Batra U, Ray S, Prabhakar SR. The influence of nickel and copper on the austempered ductile iron. Mater Eng Perform 2003;12(4):426–9. [11] Batra U, Ray S, Prabhakar SR. Effect of austenization on austempering of copper alloyed ductile iron. Mater Eng Perform 2003;12(5):597–601. [12] Bayati H, Elliott R. Austempering process in high manganese alloyed cast iron. Mater Sci Technol 1995;11:118–29. [13] El-Kashif E, El-Banna E, Riad S. Stepped austempering of GGG40 ductile cast iron. ISIJ Int 2003;43(7):1056–62.
Microstructure and fracture of alloyed austempered ductile iron Olivera Erić a , Dragan Rajnović b , Slavica Zec a , Leposava Sidjanin b , Milan T. Jovanović a,⁎ Institute of Nuclear Sciences “Vinča”, Belgrade, Serbia and Montenegro University of Novi Sad, Faculty of Technical Sciences, Novi Sad, Serbia and Montenegro a
b
Received 13 July 2005; accepted 6 January 2006
Abstract An investigation has been conducted on two austempered ductile irons alloyed with Cu and Cu + Ni, austenitized at 900°C and austempered at 350°C. The microstructure and fracture mode developed through these treatments have been identified by means of light and scanning electron microscopy and X-ray diffraction analysis. Impact energy measurements were performed on un-notched Charpy specimens. The maximum value of retained austenite volume fraction observed in the material alloyed with Cu + Ni was higher than in that alloyed with Cu austenitized and austempered under the same conditions. This led to the material alloyed with Cu + Ni having higher impact energy and substantial plastic deformation. © 2006 Elsevier Inc. All rights reserved. Keywords: Alloyed ductile iron; Bainitic transformation; Retained austenite; Impact energy
1. Introduction In recent years, there has been a significant interest in the processing and development of austempered ductile cast irons (ADI). Heat-treated ADI have been reported to have excellent mechanical properties and it appears that this material can be processed into major engineering components with a wide range of versatile properties [1–5]. ADI is a heat-treated alloyed nodular cast iron with a microstructure that consists of high carbon austenite and acicular bainitic ferrite with graphite nodules dispersed in the matrix. When discussing the microstructure of ADI, it is necessary, according to some authors [6], to ⁎ Corresponding author. Tel.: +381 11 458 222; fax: +381 11 3440 100. E-mail address: [email protected] (M.T. Jovanović). 1044-5803/$ - see front matter © 2006 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2006.01.014
distinguish the residual austenite that exists at the isothermal transformation temperature from the retained austenite that remains untransformed at ambient temperature. During the isothermal transformation, carbon partitions between ferrite laths and austenite. Undesirable phases, such as martensite and iron carbides, may also be present in smaller quantities, but it is understood that a maximum retained austenite volume fraction is important if good mechanical properties are to be achieved. The base iron chemistry and alloy additions to ductile iron play important roles in ADI technology. The addition of alloying elements during the production of ADI is often considerably higher than the levels used in the production of conventional grades of ductile irons [7]. The alloy in the austenitized condition (obtained by holding at a temperature in the range 820–950 °C long enough to dissolve the austenite with up to 2.1 wt.% carbon) is quenched to the austempering temperature
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and subjected to isothermal transformation at temperatures between 250 and 450°C before cooling to ambient temperature. The microstructure can be controlled through the variation of alloying elements, temperature and the time of austempering. The choice of austempering parameters is critical. The isothermal heat treatment should be terminated at the point that allows an optimum balance between bainitic ferrite and retained austenite, but before carbide precipitation. The formation of carbides may cause the structure to become embrittled and thus degrade the mechanical properties of ADI [5]. Since the majority of ADI components should possess satisfactory austemperability, alloying elements serve to delay the transformation of austenite [8–13]. Due to a severe restriction on the maximum thickness of bulk unalloyed ADI that can be heat-treated successfully, alloying with Si, Mn, Mo, Ni and Cu is aimed at providing a desired hardening ability. The importance of alloying additions depends on their relative effectiveness on the austempering process. Si has a very important role because the addition of more than 2 wt. % Si may retard carbide precipitation, producing carbide-free austenite in a bainitic microstructure. The effect of Si and of individual or combined additions of Cu, Mo, Ni, and Mn on the transformation characteristics of ADI have been reported earlier [9,10]. Mn and Mo delay the austempering process, but Cu does not affect the carbon diffusion in austenite or the stability of austenite. However, it has been reported that Cu suppresses carbide formation in lower bainite, whereas Ni, which acts in a similar way to Cu if present in excess of 0.5%, slows down the bainitic reaction, causing the formation of martensite at austenite cell boundaries on cooling [10]. The focus of this paper is the influence of additions of Cu and Cu + Ni on the properties of ADI. The specific items of study are the effects of austempering time and alloy composition on the microstructure, the impact properties and the fracture mechanism after austenitizing at 900°C and subsequent austempering at 350°C.
spectrometer, was as follows: (a) 3.6 C; 2.5 Si; 0.28Mn; 0.04Cr; 0.45Cu; 0.01P; 0.014 S; 0.066 Mg, and (b) 3.07C; 2.15.Si; 0.26Mn; 0.04 Cr; 1.6Cu; 1.5 Ni (the balance in both alloys was Fe). After austenitizing in a protective argon atmosphere at 900 °C for 2h, the specimens were rapidly transferred to a salt bath at the austempering temperature of 350 °C, held there for between 1 and 6 h, then air-cooled to room temperature. Fracture toughness was measured on un-notched Charpy specimens (55 mm × 10 mm × 10 mm) machined from the Y blocks. The austempered Charpy specimens were tested at room temperature in a standard impact testing machine using an applied load of 150 J. At least three specimens were tested for each heat treatment. The fracture surfaces were examined by a JEOL JSM6460LV scanning electron microscope (SEM) operating at 25 kV. Standard metallographic preparation techniques (mechanical grinding and polishing followed by etching in Nital) were applied prior to light microscope (LM) examinations. A Leitz metallographic microscope was used for microstructural characterization, and an
2. Experimental details Two ductile irons were produced in a commercial high frequency electro-induction foundry furnace. The Fe–Si–Mg (40 wt.%, 50 wt.% and 5 wt.%, respectively) alloy was used for the Mg treatment of the molten metal. The inoculant Fe–Si (with 80 wt.% Si) was added to the molten metal immediately before casting. The melt was poured from about 1420 °C into standard 25.4 mm Y block sand molds (ASTM A-395). The chemical composition (in wt.%) of these irons, obtained by
Fig. 1. LM. Microstructure of ADI alloyed with (a) Cu and (b) Cu + Ni. GN—graphite nodule; F—ferrite; P—pearlite.
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analyzed in order to calculate the volume fraction of phases present, i.e.: Iγ =Iα ¼ Rγ =Rα d Xγ =Xα where Iα and Iγ are integrated intensities from a given hkl of α and γ phase, Xα and Xγ are the volume fractions
Fig. 2. LM. Microstructure of ADI alloyed with Cu after austempering at 350°C for (a) 1h, (b) 2h, and (c) 6h. RA—retained austenite; AF— acicular ferrite; M—martensite; AFS—acicular ferrite sheaves.
Opton Axioplan light microscope equipped with the software “Vidas” program was applied to the measurement of the distribution of graphite nodules. Changes in the volume fraction of austenite retained during austempering were determined by the X-ray diffraction technique using a Siemens D-500 diffractometer with Ni-filtered Cu Kα radiation. The diffractograms were
Fig. 3. LM. Microstructure of ADI alloyed with Cu + Ni after austempering at 350°C for (a) 1h, (b) 3h and (c) 6h. RA—retained austenite; AF—acicular ferrite; M—martensite; AFS—acicular ferrite sheaves; MB—mixed bainite; FF—feathery bainitic ferrite.
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of α and γ phase. Rα and Rγ values may be calculated according to the equation
diffraction analysis was also used for identification of the phases present.
R ¼ ð1=m2 Þ½jF 2 jpð1 þ cos2 2hÞ=sin2 hcoshÞe−2M
3. Results and discussion
where F is a structural factor, p multiplicity factor, θ the Bragg angle, ν the volume of the unit cell, (1 + cos22θ) / sin2θcosθ) corresponds to Lorenz polarization factor, and e− 2M is a temperature factor (a function of θ). X-ray
3.1. Microstructure Light micrographs of cast ADI alloyed with Cu and with Cu + Ni are shown in Fig. 1(a) and (b), respectively.
Fig. 4. X-ray diffraction patterns of ADI alloyed with (a) Cu and (b) Cu + Ni, after austempering at 350°C.
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Fig. 5. The variation of retained austenite content with austempering time at 350°C in ADI alloyed with Cu and Cu + Ni.
In the microstructure of the ADI alloyed with Cu, the pearlite, at 70vol.%, outweighs the ferrite (30 vol.%) and the mainly spherical graphite nodules (the nodule count was around 100 nodules/mm2) (Fig. 1(a)). The microstructure of the ADI alloyed with Cu + Ni consists of graphite spheroid nodules (nodule count over 100 nodules/mm2) in a fully pearlitic matrix (Fig. 1(b)). The microstructures of the ADI alloyed with Cu and austempered at 350°C for 1, 2 and 6 h are shown in Fig. 2(a)–(c), respectively. In specimens austempered for 1h the predominant phase is martensite, whereas only small amounts of acicular ferrite (lower bainite), sporadically packed in sheaves, and retained austenite are present (Fig. 2(a)). At short austempering times, the carbon content is insufficient to retain stable austenite, which consequently transforms to martensite during air-cooling to the room temperature. However, after austempering for 2h, the presence of martensite could not be detected and the structure consisted only of acicular bainitic ferrite and retained austenite (Fig. 2(b)). This indicates that carbon enrichment was sufficient to stabilize the austenite even after cooling. During longer austempering (6h), the amount of retained austenite decreases and acicular bainitic ferrite dominates the microstructure. A small fraction of martensite is also present (Fig. 2(c)). In this case, the large amount of bainitic ferrite leaves less austenite available for retention. Also, the formation of bainitic ferrite consumes carbon, rendering the austenite unstable and promoting transformation to martensite. Microstructures of ADI alloyed with Cu + Ni austempered at 350°C for 1, 3 and 6 h are shown in Fig. 3 (a)–(c), respectively. After austempering for 1 h, the microstructure reveals blocky darkened areas containing
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a high fraction of plate-like martensite with a low fraction of austenite (Fig. 3(a)). After 3 h of austempering, there is an appreciable decrease in the amount of martensite and an increase in the amount of acicular bainitic ferrite and retained blocky austenite (Fig. 3(b)), suggesting that a balance between volume fractions of acicular bainite and stable retained austenite might be reached. After austempering for 6 h, only small amounts of martensite and retained austenite are present in the microstructure, and a mixture of acicular and feathery bainitic ferrite (looks like upper bainite) prevails (Fig. 3 (c)). Similar to ADI alloyed with Cu, a decrease in the amount of retained austenite was the consequence of a high volume fraction of acicular and plate bainitic ferrite. Also, there is some evidence of a mixed mode of transformation, i.e. in some areas acicular and plate bainitic ferrite overlap, a phenomenon already reported in the literature [13]. Comparing the austempered microstructures of both ductile irons, one could observe that alloying with Cu + Ni promotes the transformation of upper bainitic ferrite. The LM observations are supported by the X-ray diffraction patterns of ADI alloyed with Cu and Cu + Ni for different austempering times, shown in Fig. 4(a) and (b), respectively. The profile fitting of the 111 and 110 lines of austenite and ferrite, respectively, were identified in nearly all cases. However, doublets, due to the tetragonality of the martensite unit cell, are superimposed on the strong 110 ferrite line and could not be resolved into separate lines because of large broadening. Diffractograms of ADI alloyed with Cu (Fig. 4(a)) show that the intensity maximum of the 111 line was reached after austempering for 2h, whereas the presence of ferrite and martensite dominate after 1 and
Fig. 6. The effect of austempering time on impact energy after austempering at 350 °C of ADI alloyed with Cu and Cu + Ni.
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6h. The situation is similar with ADI alloyed with Cu + Ni, except that the 111 line could not be detected after the shortest austempering time. The presence of a broad 110 line suggests that ferrite and martensite are the main microconstituents (Fig. 4(b)). After austempering for 3 h, the 111 line exhibits an intensity maximum, indicating a high volume fraction of austenite in the microstructure. Fig. 5 shows the variation of retained austenite content with austempering time in both ductile irons. These results show that the trend in variation of the volume fraction of retained austenite with austempering time is similar for both materials. The maximum value of retained austenite (16.5 vol.%) noted in the ADI alloyed with Cu is achieved after 2h. However, with the ADI alloyed with Cu + Ni, the retained austenite reaches a maximum (19 vol.%) after 3 h. It appears that alloying with Cu + Ni delays the transformation kinetics of residual austenite (isothermal) until the maximum of
the volume fraction is reached. These results are in agreement with the LM and X-ray diffraction findings. 3.2. Impact properties The effect of alloying ADI with Cu and Cu + Ni on the impact energy of specimens previously austempered at 350 °C is illustrated in Fig. 6. The results show that the impact energy increases as the austempering time increases up to a maximum value, after which it decreases. The maximum impact energy (110 J) of ADI alloyed with Cu is achieved after austempering for 2 h, whereas alloying with Cu + Ni postpones the maximum (122 J in this case) to 3 h. The fact that the maximum impact energy corresponds to the maximum value of retained austenite in both cases suggests that the impact values depend on the retained austenite content. That is, with increasing content of retained austenite up to the maximum value, the impact energy also increases,
Fig. 7. SEM. Fracture morphology after austempering at 350 °C. ADI alloyed with Cu for (a) 2h and (c) 6h; ADI alloyed with Cu + Ni for (b) 3 h and (d) 6h.
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then, as the retained austenite content decreases, so the impact energy starts to decrease.
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of retained austenite up to a maximum value impact energy increases, then a decrease occurs with the decrease of retained austenite.
3.3. Fractography Acknowledgements The fracture surfaces of un-notched impact specimens after austempering at 350 °C for times corresponding to maximum impact energy (2 and 3h) and to the subsequent lowest impact energy (after 6h) are shown in Fig. 7. The micrographs of specimens with highest values of retained austenite volume fraction and impact energy exhibit a typical ductile fracture mode (Fig. 7(a) and (b)). The somewhat higher number of dimples observed in the Cu + Ni alloy (Fig. 7(b)) is associated with the higher values of toughness and retained austenite volume fraction. After austempering for 6h, the fracture mode is changed. In the ADI alloyed with Cu a mixture of ductile and cleavage fracture prevails (Fig. 7(c)) but in the Cu + Ni alloyed ADI, only cleavage fracture can be seen (Fig. 7(d)). These results are in accordance with the results of the impact energy testing because the Cu alloyed ADI exhibited the higher impact energy after austempering for 6 h. 4. Conclusions The effect of alloying with Cu and Cu + Ni on the microstructure and impact properties of austempered ductile iron has been studied. Addition of Cu + Ni delays the transformation kinetics of residual austenite (isothermal) resulting in a shift of the maximum of volume fraction of retained austenite to 3 h of austempering, compared to 2 h in ADI alloyed with Cu. In the same time, higher maximum value of the volume fraction of retained austenite was observed in ADI alloyed with Cu + Ni. In the same ADI a substantial plastic deformation at the peak of impact energy is associated with the highest volume fraction of retained austenite. It has been demonstrated that the volume fraction of retained austenite strongly affects impact energy of both irons, i.e. with increasing content
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